Toughening mechanisms in cellulose nanopaper: the contribution of amorphous regions

Cellulose nanopaper is a strong and tough fibrous network composed of hydrogen bonded cellulose nanofibres. Upon loading, cellulose nanopaper exhibits a long inelastic portion of the stress–strain curve which imparts high toughness into the material. Toughening mechanisms in cellulose nanopaper have been studied in the past but mechanisms proposed were often rather speculative. In this paper, we aim to study potential toughening mechanisms in a systematic manner at multiple hierarchical levels in cellulose nanopaper. It was proposed that the toughness of cellulose nanopaper is not, as is often assumed, entirely caused by large scale inter-fibre slippage and reorientation of cellulose nanofibres. Here it is suggested that dominant toughening mechanism in cellulose nanopaper is associated with segmental motion of molecules facilitated by the breakage of hydrogen bonds within amorphous regions .

Toughening mechanisms in cellulose nanopaper: the contribution of amorphous regions

Toughening mechanisms in cellulose nanopaper: the contribution of amorphous regions
Rui Mao . Nan Meng . Wei Tu . Ton Peijs 0 1
0 W. Tu Nanoforce Technology Ltd., Joseph Priestley Building, Queen Mary University of London , Mile End Road, London E1 4NS , UK
1 R. Mao N. Meng T. Peijs (&) School of Engineering and Materials Science, Queen Mary University of London , Mile End Road, London E1 4NS , UK
Cellulose nanopaper is a strong and tough fibrous network composed of hydrogen bonded cellulose nanofibres. Upon loading, cellulose nanopaper exhibits a long inelastic portion of the stress-strain curve which imparts high toughness into the material. Toughening mechanisms in cellulose nanopaper have been studied in the past but mechanisms proposed were often rather speculative. In this paper, we aim to study potential toughening mechanisms in a systematic manner at multiple hierarchical levels in cellulose nanopaper. It was proposed that the toughness of cellulose nanopaper is not, as is often assumed, entirely caused by large scale inter-fibre slippage and reorientation of cellulose nanofibres. Here it is suggested that dominant toughening mechanism in cellulose nanopaper is associated with segmental motion of molecules facilitated by the breakage of hydrogen bonds within amorphous regions .
Nanocellulose; Nanopaper; Yielding; Plasticity; Toughness
Introduction
Cellulose widely exists in plant cell walls in the form
of microfibrils
(Barnett and Bonham 2004; Keckes
et al. 2003)
. The repeat unit in the cellulose molecule
consists of two anhydroglucose rings which are linked
by b-1,4 glycosidic bonding (Klemm et al. 1998).
Cellulose nanofibres can be extracted in the form of
individual microfibrils and/or their aggregations by
deconstructing the cell wall structure
(Eichhorn et al.
2010)
. Cellulose nanofibre is composed of crystalline
regions where cellulose chains are believed to align
parallel to the axis of the nanofibre together with
amorphous regions where cellulose chains are
notordered
(Haslach 2000)
. Crystalline regions are
separated by amorphous regions at intervals of up to
800 A˚ , with individual cellulose molecules
connecting several crystalline and amorphous regions
(Haslach 2000; Salme´n 1986)
. The amorphous
cellulose is also present at the surface of the crystallites.
Cellulose crystals prepared via acid hydrolysis
displayed a crystallinity index of about 80%
(Kargarzadeh et al. 2012)
, which suggests a contribution
of cellulose chains within the crystallite surface to the
amorphous phase of the nanofibres. Cellulose
nanopaper is a fibrous network composed of cellulose
nanofibres connected together by van der Waals forces
and hydrogen bonding. Because of improved
mechanical properties of nanofibres over micro-sized
cellulose fibres, in combination with an increased network
density, it shows superior mechanical properties
compared with conventional paper. Tensile strengths
and Young’s moduli of cellulose nanopaper have been
reported to be around 200 MPa and 13 GPa,
respectively
(Henriksson et al. 2008)
. Its strain-to-failure can
reach a value as high as 10%, imparting high
toughness into the cellulose nanopaper (work to
fracture = 15.1 MJ m-3)
(Henriksson et al. 2008)
.
Typical stress–strain curves of cellulose nanopaper
show an elastic region followed by a long inelastic
region where the toughness primarily originates from.
Therefore, the exploration of toughening mechanisms
in cellulose nanopaper should focus on mechanisms of
inelasticity in such nanopapers.
So far, many papers have focused on the structure–
property relation of cellulose nanopaper, while the
reported mechanisms are all rather speculative.
Hsieh
et al. (2008)
proposed that the inelastic portion in the
stress–strain curve is related to breakdown of the
fibrous network by bond breaking and fibre pull-out.
They also proposed that twisting of individual
cellulose fibrils could possibly contribute to inelasticity.
Henriksson et al. (2008)
prepared cellulose
nanopapers with different porosities by drying cellulose
nanopapers from solvents of various polarities. They
speculated that inelasticity is related to inter-fibrillar
debonding and slippage of nanofibrils promoted by
voids.
Sehaqui et al. (2011)
reported that higher
porosity and higher surface area led to weaker
interfibre bonds, resulting in a lower tensile modulus and
higher strain-at-break of nanopapers and a weaker
strain hardening behaviour in the inelastic region of
the stress–strain curve. The dependence of yield stress
on the degree of polymerization of the cellulose
molecules and length of the nanofibres was also
investigated but no clear trend was observed
(Fukuzumi et al. 2013; Henriksson et al. 2008)
. The inelastic
behaviour can be changed by changing humidity
(Benitez et al. 2013). It was reported that inelasticity is
facilitated by high humidity, and it was proposed that
here inelasticity is caused by inter-fibrillar debonding
and possible sliding. Inter-fibrillar hydrogen bonds
can be weakened and broken by water molecules,
leading to breakage of hydrogen bonds and reduced
inter-fibrillar friction
(Benitez et al. 2013; Quero et al.
2011)
. Robust cellulose nanopapers have been
produced by hot pressing, with inelastic behaviour being
more pronounced for less well pressed nanopapers
(O¨ sterberg et al. 2013). It was speculated that
nanopapers that were hot pressed for longer periods
of time were denser and had fewer voids. By this more
hydrogen bonds could form after extended periods of
high pressure and temperature. The effect of
microvoids on the mechanical properties of all-cellulose
fibreboard was also reported by
Are´valo and Peijs
(2016)
where denser fibreboard resulted in higher
mechanical properties. Recently,
Zhu et al. (2015)
proposed that the high toughness of cellulose
nanopaper was caused by breaking and reformation of
hydrogen bonding during inter-fibre slippage. The
hypothesis was supported by results of atomistic
simulation. Structure–property relationships of
cellulose nanopaper were also studied using finite element
modelling of fibrous networks
(Mao et al. 2017)
. From
literatures it can be concluded that the inelasticity of
cellulose nanopaper is generally speculated to be
associated with the break-down of hydrogen bonds
between cellulose nanofibres together with inter-fibre
slippage. However, so far no direct evidence exists for
such a hypothesis.
Cellulose nanofibres are composed of
macromolecules, which are organized in crystalline and
amorphous domains or regions. Therefore there is a
distinct possibility that the non-linear behaviour in
these materials is related to the same phenomena that
promote inelastic behaviour in semi-crystalline
polymers. This paper will therefore focus on mechanisms
for inelasticity in cellulose nanopaper at different
length scales. First, a two-dimensional digital image
correlation (2D DIC) technique, which has been used
earlier to characterize the strain distribution in
paperboard by
Hagman and Nyga˚rds (2012)
was applied to
cellulose nanopaper. Next, the possibility of inter-fibre
slippage between individual cellulose nanofibres and
subsequent reorientation of nanofibres during the
inelastic stage was examined using repeated
loading–unloading experiments, in situ Raman
spectroscopy and wide angle X-ray diffraction (WAXD).
Finally, the relationship between local mobility of
molecular chain segments in amorphous domains and
inelastic behaviour was investigated using polarized
optical microscopy (POM), tensile testing performed
at different strain rates and temperatures, as well as
temperature dependent dielectric spectroscopy.
Materials and methods
Materials
Preparation of cellulose nanofibres
Nanofibrillated cellulose (NFC) nanofibres were
kindly provided by Prof. Berglund’s group at KTH
Stockholm (Sweden). The NFC was prepared from
softwood pulp provided by Nordic Pulp and Paper
(Sweden). The lignin and hemicellulose contents were
0.7 and 13.8%, respectively. Following
Henriksson
et al. (2007)
, the pulp was dispersed in deionized water
followed by a pretreatment process involving a
combination of an enzymatic treatment (Novozym
476, Novozymes A/S, Denmark) and mechanical
beating in a laboratory beater (PFI mill, Hamjern
Maskin, Norway). This pretreated pulp was
homogenized by 8 passes through a microfluidizer (M-110EH,
Microfluidics Inc., USA). From this a water
suspension with a solid concentration of 1.7% NFC was
obtained.
Preparation of cellulose nanopaper
Cellulose nanopaper was prepared by first mixing
NFC suspension with 0.2 g dry weight and
approximately 40 mL distilled water. The speed of the
mixer (IKA Ultra Turrax mixer D125 Basic,
Germany) was 12,000 rpm and the mixing time was
10 min. Then the diluted suspension was degassed.
Finally, the suspension was poured into a petri-dish
followed by conditioning in an incubator at 37 C
for several days. During this time, silica-gel
particles were also placed in the incubator in order to
absorb the water vapour. The resulting cellulose
nanopaper had a thickness of approximately 30 lm.
The grammage and porosity of the cellulose
nanopaper were 46 g m-2 and 11%, respectively.
The porosity was calculated using Eq. 1.
Porosity ¼ 1
qnanopaper
qcellulose
where qnanopaper is the density of cellulose nanopaper
and qcellulose is the density of cellulose which is
1500 kg m-3
(Henriksson et al. 2008)
.
ð1Þ
Characterizations
Tensile testing
Tensile tests were conducted using universal testing
machines (Instron 5566 and Instron 5584, USA). The
load cell for all the tensile tests was 100 N. Specimens
of 5 mm in width were used and conditioned at a
temperature of 23 C and a relative humidity (RH) of
50% for at least 24 h before testing. The distance
between the clamps was 20 mm. First, tensile tests were
performed at a speed of 2 mm min-1 using an Instron
5566 to obtain the stress–strain curve and mechanical
properties of the nanopapers. Then, tensile tests were
conducted at different strain rates and temperatures
using an Instron 5584 equipped with an environmental
chamber for temperature control. Environmental
conditions during testing included: (1) a temperature of
23 C and RH of 50% and (2) a temperature of 50 C
and RH of 50%, respectively. The results were based on
at least six specimens for each sample. Strain was
calculated using the cross-head displacement since the
influence of frame compliance (\0.1%) can be
neglected for the loads involved in these tests.
Repeated loading–unloading testing
Repeated loading–unloading experiments were
performed on cellulose nanopaper to extract Young’s
modulus and yield stress for each cycle. The tests were
carried out using an Instron 5566 universal testing
machine (USA) equipped with a 100 N load cell.
Specimen dimensions and conditioning (23 C and
50% RH) were the same as in static tensile testing. The
loading speed was 2 mm min-1. The specimen was
first loaded in the elastic region to obtain the elastic
modulus. Then it was repeatedly loaded to different
peak loads in the inelastic region where peak load was
set to increase incrementally for each load cycle. The
test was performed in triplicate to obtain averages.
2D Digital image correlation (2D DIC)
2D DIC technique was used to map the strain
distribution in the tensile specimens. Specimen
dimensions and conditioning were the same as in
tensile testing. The surface of the specimens was first
spray coated with white paint to obtain a clear
background (*6 lm in thickness). Then, black
speckles of *100 lm were sprayed on the surface
of the specimen to achieve good contrast. The
specimens were loaded in a micro-tensile tester
(Deben, UK) equipped with a 200 N load cell.
Distance between the clamps was 10 mm and the test
speed was set at 0.2 mm min-1. Images were taken by
a CCD camera during loading. The pixel size for each
picture was 9.0 lm 9 9.0 lm. Displacement between
speckles during loading was analysed by Ncorr, a
Matlab programme developed at Georgia Institute of
Technology, USA
(Blaber et al. 2015)
.
Wide angle X-ray diffraction (WAXD)
Crystallinity of cellulose nanofibres forming cellulose
nanopaper was determined from one-dimensional
wide-angle X-ray diffraction (1D WAXD) patterns in
a 2h range of 5 –70 . The patterns were obtained using
a Bragg–Brentano geometry X-ray diffractometer
(X’Pert Pro, PANalytical, The Netherlands) equipped
with Cu/Ka radiation. The obtained spectra were
composed of sharp peaks from crystallites and broad
background scattering from amorphous regions.
Crystallinity of cellulose nanofibres forming cellulose
nanopaper was calculated using Eq. 2 where the
crystalline (P Ac) and amorphous fractions (P Aa)
were obtained by fitting the spectra into 4 crystalline
peaks and a broad band at approximately 21.5 which
was assigned to amorphous contributions according to
Park et al. (2010)
. The peak fitting was performed
using PeakFit software (www.systat.com).
P Ac
vc ¼ P Ac þ P Aa
100%
ð2Þ
The crystalline preferred orientation of cellulose
nanofibres was examined using two-dimensional
wideangle X-ray diffraction (2D WAXD) ring patterns,
which were obtained using a transmission geometry on
a single crystal X-ray diffractometer (Kappa ApexII
Duo, Bruker AXS GmbH, Germany). The
measurements were carried out with the direction of the X-ray
beam perpendicular to the plane of nanopaper,
resulting in patterns in the in-plane direction. The 2D
WAXD measurement was conducted on specimens
before straining and post-fracture. Before straining, the
X-ray was diverted to the specimen surface, while for
failed specimens the X-ray was diverted to an area
close to the fracture site. The surface area covered by
2
X-ray was approximately 0.00785 mm .
Polarized optical microscopy (POM)
POM was used to examine the degree of anisotropy of
specimens before and after tensile testing. The
examination of molecular anisotropy is based on the fact
that a linearly polarized incident light beam travelling
through a uniaxially oriented material is split into two
orthogonally polarized rays (ordinary and
extraordinary rays) with different phase velocities so that a
phase difference (d) can be identified between the rays
emerging from the material
(Bhupathi et al. 2010)
. If a
light beam with intensity of I0 is incident on an
anisotropic material which is placed between two
crossed polarizers, the transmitted light intensity I can
be given by Eq. 3 (Born and Wolf 1999).
ð3Þ
I ¼ I0 sin 2u sin2 d
2
where u is the angle between the optical axis and the
polarization axis of the polarizer. In our case, the
optical axis was assumed to be along the straining
direction. For an isotropic material, the optical image
will be dark no matter which direction the specimen is
oriented. For an anisotropic material, the minimum
transmitted light intensity (dark) can be obtained when
the straining direction is along the direction of the
polarizers. Maximum intensity (bright) can be
obtained when the angle between straining direction
and the polarization direction of each polarizer is 45 .
In the POM measurements, specimen dimension
and conditioning were the same as in tensile testing. A
specimen was first placed between two crossed
polarizers where the direction of tension was aligned
with one polarization direction (0 ). One image was
taken at a magnification of 59. Another image was
also taken when the same specimen was placed at an
angle of 45 . Then, the specimen was stretched until
failure using an Instron 5566 (USA) equipped with a
100 N load cell at a crosshead speed of 2 mm min-1.
Images of the failed specimen were taken close to the
location of fracture, with the specimen being placed at
angles of 0 and 45 relative to the polarization axis,
respectively. A schematic of the test configuration is
shown in Fig. 1. Anisotropy was identified if a
specimen placed at an angle of 45 showed a brighter
image than at 0 .
Fig. 1 Schematic of the
experimental test setup for
POM
In-situ Raman spectroscopy
The relationship between Raman band shifts and
applied stresses and strains was investigated by
recording Raman spectra using a Renishaw Raman
Spectrometer (UK) equipped with a 633 nm Helium–
Neon laser with a power of 35 mW during tensile
testing. The laser power at specimen surface was
1–2 mW. The 1800 lines mm-1 diffraction grating
was used and the resolution was 1 cm-1. Specimen
dimensions and conditioning were the same as in
tensile testing. A specimen was mounted on a
microtensile tester (Deben, UK) equipped with a 200 N load
cell and loaded at a speed of 0.2 mm min-1. The
specimens were strained step-wise (0.2% strain
increment), and at each step, the specimen was excited by a
laser beam. A 509 objective lens was used to focus the
laser beam on the surface of the specimen, while the
spot size of the laser was *2 lm. The wavenumbers
used in the Raman spectra varied from 1050 and
1150 cm-1. The Raman spectra were recorded using
an exposure time of 10 s and three accumulations. The
peak initially positioned at approximately 1095 cm-1
was determined by fitting the Raman band using a
Lorentzian function. This test was triplicated.
Temperature dependent dielectric spectroscopy
The temperature dependence of the dielectric loss was
measured using a LCR meter (4284A, Agilent
Technologies, USA) connected to an in-house made oven.
The measurements were conducted in a temperature
range of -180 to 245 C at four different frequencies,
100, 10 kHz, 1000 and 100 Hz.
Results and discussion
Fig. 2 Stress–strain curve of cellulose nanopaper. The red line
sections the stress–strain curve into elastic and inelastic regions
based on the 0.2% offset ‘yield’ point. The insets are 2D DIC
strain distribution plots at different levels of strain (scale bar
1 mm). The degree of strain is indicated by the far right bar in
the graph. (Color figure online)
Work to fracture (MJ m-3)
approximately 1%, after which a long inelastic region
occurs until final failure. The strain distribution in the
nanopaper as measured by 2D DIC at different loading
stages are shown as insets. The strain distribution
image in the elastic region shows a relatively uniform
strain distribution in the nanopaper plane, while a less
uniform strain distribution exists in the inelastic
region. At the latter stage, some nanofibres are less
loaded than others with strain values at every position
in the nanopaper exceeding the value at the transition
from elastic to inelastic region, indicating that
inelasticity occurs throughout the whole cellulose
nanopaper.
Repeated loading–unloading experiments were
carried out to investigate the effect of inelastic loading
history on Young’s modulus and yield stress (Fig. 3a).
Here, the words ‘plastic deformation’ and ‘yield
stress’ are used only based on phenomenological
interpretations since the mechanisms of inelastic
deformation remain unclear. The development of
Young’s modulus with number of cycles is shown in
Fig. 3b. The specimens were first strained in the
elastic region (loading cycle 1 and 2) in order to obtain
the initial Young’s modulus. The Young’s modulus at
cycle 1 and 2 is 9.6 ± 0.7 GPa. Subsequent loading
cycles were performed in the inelastic region and it is
shown that the Young’s modulus remains constant for
all cycles. Furthermore, the values of Young’s
modulus are similar in both the elastic and inelastic
regions. This result is consistent with the data reported
by
Henriksson et al. (2008)
who measured the
Young’s modulus of cellulose nanopaper as a function
of the number of loading–unloading cycles, and
reported that the modulus during loading remained
relatively unchanged with cycles. The values of yield
stress of the cellulose nanopaper against loading
cycles are plotted in Fig. 3c. Also yield stress is fairly
independent of number of loading cycles and does not
show any clear trend. The results of these repeated
loading–unloading tests indicate that inelastic
deformation has no apparent effect on Young’s modulus
and yield stress. Therefore, the occurrence of large
scale nanofibre reorientation in the inelastic region is
not very plausible as this would affect these properties.
Raman spectroscopy was used to examine the
micromechanics of cellulose nanopaper in both the
elastic and inelastic region. Figure 4a shows the
Raman spectra of cellulose nanopaper at different
levels of strain where a band initially at approximately
1095 cm-1 which corresponds to the C–O and C–C
stretching mode in the cellulose backbone can be
observed for an undeformed specimen
(Gierlinger
et al. 2006; Quero et al. 2010; Wiley and Atalla 1987)
.
This Raman band shifted to lower wavenumber
positions when the cellulose nanopaper was stretched.
This shift can be ascribed to molecular straining of the
cellulose molecule
(Eichhorn et al. 2001a)
. The shift
of the Raman peak positions are plotted against
applied stress and strain in Fig. 4b, c, respectively
with the data fitted by linear functions with high
correlation. The drawn lines in Fig. 4b, c are lines with
gradient values that correspond to the ones listed in the
figures. This indicates that cellulose molecules are
stressed proportionally with the applied stress and/or
strain. As the Raman band shift reflects the extent of
molecular straining, the band shift rate (slope of
Raman band shift against stress) will change if the
contribution of the applied stress to molecular
deformation changes
(Eichhorn et al. 2001b)
. Therefore, it
can be concluded that no change in such contribution
was detected by Raman spectra.
One can imagine that if the inelasticity is dominated
by large scale inter-fibre slippage and reorientation,
the contribution of the applied stress to molecular
straining will change when inelasticity takes place
since slippage and reorientation in fibre scale would
cause massive unloading in molecular chains.
However, Raman spectra show that such a change cannot
be detected, meaning that inelasticity is not dominated
by large scale inter-fibre slippage and reorientation.
The origin of inelasticity is therefore likely to be
related to processes at a lower hierarchical level. The
Raman band shift rate in Tanpichai et al’s (2012) paper
was -0.5 cm-1 %-1 where its absolute value is
Fig. 4 a Raman band shift of 1095 cm-1 peak against strain,
b the Raman band shift in the 1095 cm-1 as a function of stress
and c the Raman band shift in the 1095 cm-1 as a function of
strain. b, c the dots of different shapes and colours represent
measurements performed on different nanopaper specimens.
The gradients in a, b are average values of slopes obtained by
linear fitting data from each specimen. The drawn lines in b,
c are lines with gradient values that correspond to the ones listed
in the figures
higher than the one measured in our current study
(-0.23 cm-1 %-1). Here, cellulose nanopaper was
prepared through a suspension casting method where
the resulting nanofibres are expected to be bend and
the number of fibre–fibre bonds is relatively low.
Tanpichai et al. prepared nanopapers by filtration of
NFC suspensions followed by constrained drying of
the filtered samples, resulting in straighter nanofibres
and more inter-fibre bonds. Hence, the stress transfer
ability in Tanpichai et al’s nanopaper is expected to be
better than in the nanopapers prepared by a suspension
casting method, resulting in a stronger Raman band
shift. In addition, the Raman band shift here cannot be
used to calculate Young’s modulus of cellulose
nanopaper since the used Raman laser beam was not
polarized. According to
Tanpichai et al. (2012)
,
a polarized laser beam is required to characterize the
Young’s modulus since the Raman band shift rate
changes with the angle between polarization direction
of the beam and strain direction. Therefore, our Raman
band shift results can only be used for qualitative
analysis.
In order to determine the crystallinity of the
cellulose nanofibres in the nanopaper, the X-ray
diffraction spectrum of cellulose nanopaper was
analysed. Figure 5 shows the X-ray diffraction pattern
of cellulose. It can be seen that the diffraction peaks
are located at 2h angles of 14.8 , 16.8 , 22.6 , 34.9 ,
which are typical native cellulose peaks
(Tingaut et al.
2010)
. The calculated crystallinity is 68%, which
means that cellulose nanopaper has amorphous
Fig. 5 X-ray diffraction spectra of cellulose nanopaper,
showing typical native cellulose diffraction peaks located at 14.8 ,
16.8 , 22.6 , 34.9
domains or regions. Therefore, both the crystalline
and amorphous phase need to be examined with
respect to their contribution to the inelastic behaviour
of cellulose nanopaper. Figure 5 also shows a reduced
intensity for the cluster of reflections near the 004 peak
at 34.9 on a Cu radiation pattern compared with the
intensity for the reflections from randomly oriented
crystallites, which indicates that crystallites in the
cellulose nanopaper do show one-dimensional
orientations
(Segal et al. 1959; French and Cintro´ n 2013;
French 2014)
.
Reorientation of nanofibers has been reported for
specimens that were being drawn at a wet state prior to
drying
(Sehaqui et al. 2012)
. Here, the nanofibers
reoriented because the hydrogen bonding was affected
by the presence of water molecules which also acted as
plasticizer between nanofibers. Therefore, nanofibers
were able to move and align in the direction of an
external force. Figure 6 shows the 2D WAXD data
reflecting the in-plane degree of orientation. Tests
were conducted on nanopapers before and after
straining to failure (in a relative dry state). In-plane
orientation of the crystallites in the paper is reasonably
random before straining. Interestingly, this random
orientation remained even when the specimen was
stretched into the inelastic region and up to ultimate
failure. The crystalline regions in the cellulose
nanofibres did not reorient along the straining
direction. This confirms our earlier hypothesis that
cellulose nanofibres do not necessarily reorient in the
inelastic region.
From above observations, it can be hypothesized
that inelasticity in cellulose nanopaper is not the result
of reorientation of cellulose nanofibres or crystalline
regions in cellulose nanofibres. Hence, it is of interest
to examine more closely the contribution of the
amorphous regions in cellulose nanofibres to the
inelastic deformation behaviour.
Figure 7 shows the images taken by POM with the
nanopaper placed between crossed polarizers.
Figure 7a, b show the optical images before straining,
with the sample being positioned at the angles of 0
and 45 relative to the polarization axis, respectively.
Images for both angles appear dark, indicating an
isotropic structure of the cellulose nanopaper.
Figure 7c, d show the optical images of samples after
failure at angles of 0 and 45 relative to the
polarization axis, respectively. The image at 0 is
dark while the image at 45 appears bright. This
chain mobility increases so that molecular strain rate is
easier to match applied strain rate resulting in a lower
yield stress. Similarly, low strain rates result in a
lowering of the yield stress
(Engels et al. 2010;
Kanters 2015; Mulliken and Boyce 2006; Van Erp
et al. 2009)
. The relationship between yield stress and
temperature and applied strain rate can be described by
the Eyring equation
(Chale´at et al. 2008; Senden et al.
2012; So¨ ntjens et al. 2012)
.
e_ ¼ e_0 exp
DH
RT
sinh
V ry
RT
ð4Þ
ð5Þ
ð6Þ
ry R
T ¼ V
ry R
T ¼ V
sinh 1
exp
_
e
_
e0
DH
RT
ln
2e_
_
e0
DH
þ RT
where e_ is strain rate, ry is the yield stress, e_0 is the
constant pre-exponential factor, DH is the activation
energy, V is the activation volume, R is the gas
constant and T is the absolute temperature. Equation 4
can be expressed as following
For large yield stress, Eq. 5 can reduce to
Fig. 6 2D WAXD patterns with the X-ray beam perpendicular
to the cellulose nanopaper surface: a before straining, b after
ultimate failure. No arcs can be observed in both images, which
indicates the absence of preferential orientation
indicates that this sample exhibits some anisotropy
upon straining in the inelastic region. Furthermore, it
can be deduced that this anisotropy originates from
amorphous regions.
Yielding in semi-crystalline or amorphous
polymers such as polyethylene terephthalate (PET)
(Lim
et al. 2003)
, polycarbonate (PC)
(Mulliken and Boyce
2006)
, poly(methyl methacrylate) (PMMA)
(Mulliken
and Boyce 2006)
, polypropylene (PP)
(Alcock et al.
2007)
, poly(vinyl alcohol) (PVA)
(Govaert and Peijs
1995)
and plasticized starch blend
(Chale´at et al.
2008)
can be attributed to changes of chain
conformation in amorphous regions, which depends on chain
mobility. Yielding takes place when the molecular
plastic strain rate resulting from chain mobility
matches the applied strain rate. At high temperature,
Figure 8 shows the strain rate and temperature
dependence of the yield stress for cellulose nanopaper.
Temperature-normalized yield stress is plotted against
the logarithmic strain rate for samples tested at
different temperatures. In this paper, the 0.2% offset
yield point is used since there is no maximum stress at
yield point in the stress–strain curve. It can be seen that
the normalized yield stress increases with strain rate
but decreases with temperature. This indicates that
lower strain rates and higher temperatures facilitate
the mobility of cellulose molecules. Equation 6 was
used to fit the data of the temperature-normalized yield
stress versus logarithmic strain rate at different
temperatures. It can be seen that the normalized yield
stress increases linearly with the logarithm of strain
rate, which indicates that one thermally activated
process can be used to describe the yield behaviour
under these conditions. The values of the activation
parameters, DH and V , are shown in Fig. 8. The
activation volume is 0.8 nm3, which is defined as the
product of cross-sectional area of the moving unit and
its moving distance
(Lim et al. 2003)
. This activation
volume is of the order of activation volume values for
polymers where molecular processes dominate
failure, 0 angle and d after ultimate failure, 45 angle. The
black cross is a marker to locate the imaging region before and
after straining
inelasticity
(Lim et al. 2003)
. The activation energy
which is defined as the energy barrier for the molecular
process is 157 kJ mol-1. This value is much higher
than the dissociation energy of hydrogen bonding
which ranges from 20 to 50 kJ mol-1
(Bohidar 2015)
but lower than the dissociation energy of C–C and
C–O covalent bonds ranging from 300 to
500 kJ mol-1
(Blanksby and Ellison 2003; Chale´at
et al. 2008)
, suggesting that hydrogen bonds were
broken when cellulose nanopaper was loaded in the
inelastic region. Therefore, this study suggests that the
inelastic behaviour of cellulose nanopaper originates
from molecular mobility in amorphous regions
favoured by breakage of hydrogen bonds.
In order to study these molecular processes in more
detail, possible relaxation processes in cellulose
nanofibres were examined using dielectric
spectroscopy. Figure 9 shows the temperature dependence
of the dielectric loss at different frequencies. Two
dielectric loss peaks can be identified at a frequency of
100 Hz. For polarized polymers, this dielectric loss is
primarily from overcoming friction during changing
of chain conformation under an alternating applied
electric field. Therefore, a dielectric loss peak
represents a state when the chain mobility does not match
the alternation of the applied electric field. The
dielectric properties of cellulose have been
extensively investigated
(Jafarpour et al. 2007; Rachocki
et al. 2005; Roig et al. 2011)
. The peak of the dielectric
loss at low temperature is usually assigned to
Fig. 9 Dielectric loss versus temperature at different
frequencies for cellulose nanopaper, showing two dielectric loss peaks
at different temperatures. The peaks move towards higher
temperatures for higher frequencies
secondary relaxation while the one at higher
temperature is attributed to primary relaxation or specifically
the dielectric manifestation of the glass transition
(Roig et al. 2011)
.
It is worth noting that dielectric loss peaks move to
higher temperature for higher frequencies of applied
electric field. This behaviour is consistent with the
strain rate and temperature dependence of the yield
stress. Since the strain rate and temperature
dependence of the yield stress is associated with a single
molecular process (motion in molecules in amorphous
regions represented by yield stress as a linear function
of strain rate in Fig. 8), one may assume that the
inelastic behaviour of cellulose nanopaper is strongly
associated with segmental motion of the molecular
main chain.
Conclusions
Cellulose nanopaper is a strong and tough fibrous
network composed of hydrogen bonded cellulose
nanofibres. Upon loading, it shows significant inelastic
behaviour which results in its high toughness. Strain
mapping using 2D DIC revealed that inelasticity
occurs all over the sample. Mechanical testing and
Raman spectroscopy indicated that inelasticity is not
the result of large scale reorientation of cellulose
nanofibres in the network. 2D WAXD patterns
suggested that also molecular reorientation in
crystalline regions is not responsible for the inelastic
behaviour. On the other hand, POM images revealed a
transition from an isotropic amorphous structure to an
anisotropic structure after straining. Eyring’s theory
was used to describe the strain rate and temperature
dependence of the yield stress. This suggests that the
inelastic behaviour is associated with a single
molecular process, which is supported by the POM results
which indicate that inelasticity takes place within the
nanofibres. The activation energy of this process was
sufficient to break the van der Waals or hydrogen
bonds between the nanofibres or cellulose molecules
but not enough to break the covalent bonds within the
individual nanofibres. The temperature dependence of
the dielectric loss showed that inelasticity occurred in
a temperature range where primary relaxation
dominates. In short, the current work hints that segment
motion of molecules in amorphous regions is an
important cause for the inelastic region in the stress–
strain curve of cellulose nanopaper at a dry state.
However, other toughening mechanisms such as fibre
slippage and reorientation may become active and
contribute to the inelastic mechanical behaviour of
cellulose nanopaper at a wet state.
Acknowledgments The authors would like to thank Dr. Rory
Wilson for the technical help in XRD. The authors would also
like to acknowledge Prof. Lars Berglund from the Wallenberg
Wood Science Center at the KTH Stockholm for providing
cellulose nanofibres and stimulating discussions on toughening
mechanisms. R.M. would like to acknowledge the China
Scholarship Council (CSC) scheme for their financial support.
Compliance with ethical standards
Conflict of interest The authors declare that they have no
conflict of interest.
Open Access This article is distributed under the terms of the
Creative Commons Attribution 4.0 International License (http://
creativecommons.org/licenses/by/4.0/), which permits
unrestricted use, distribution, and reproduction in any medium,
provided you give appropriate credit to the original
author(s) and the source, provide a link to the Creative
Commons license, and indicate if changes were made.
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